Heat treatment of alloys having elements for improving grain boundary strength

ABSTRACT

Heat treatment of alloys having elements for improving grain boundary strength is disclosed. Components directly after castings often reveal low or no transverse grain boundary strength, so that cracks do appear and decrease the yield rate. The provided measures do not lead to low transverse grain boundary strength but maintains efficient grain boundary strength, so that the yield rate of components without cracks is increased.

CROSS REFERENCE TO RELATED APPLICATIONS

This application is a continuation-in-part of U.S. application Ser. No. 10/641,995, filed Aug. 15, 2003, which is the U.S. National Stage of International Application No. PCT/EP02/11856, filed Oct. 12, 2002, and claims the benefits of these applications. The International and U.S. Applications are incorporated by reference herein in their entirety.

FIELD OF THE INVENTION

The present invention relates to a heat treatment of alloys, especially nickel base superalloy and, more particularly, to castings having a columnar grain microstructure.

BACKGROUND OF THE INVENTION

U.S. Pat. No. 4,597,809 describes single crystal castings made from a nickel base superalloy having a matrix with a composition consisting essentially of, in weight %, of 9.5% to 14% Cr, 7%. to 11% Co, 1% to 2.5% Mo, 3% to 6% W, 1% to 4% Ta, 3% to 4% Al, 3% to 5% Ti, 6.5% to 8% Al+Ti, 0% to 1% Nb, and balance essentially nickel with the matrix containing about 0.4 to about 1.5 volume of a phase based an tantalum carbide as a result of the inclusion in the alloy of about 0.05% to about 0.15% C and extra Ta in an amount equal to 1 to 17 times the C content.

Single crystal castings produced from the aforementioned nickel base superalloy exhibit inadequate transverse grain boundary strength. The present inventors attempted to produce directionally solidified (DS) columnar grain castings of the nickel base superalloy. However, the directionally solidified (DS) columnar grain castings produced were unacceptable as DS castings as a result of the castings exhibiting essentially no transverse grain boundary strength and no ductility when tested at a temperature of 750 degrees C. (1382 degrees F.) and stress of 660 MPa (95.7 Ksi). The transverse grain boundary strength and ductility were so deficient as to render DS columnar grain castings produced from the aforementioned nickel base superalloy unsuitable for use as turbine blades of gas turbine engines.

WO 99/67435 discloses nickel base superalloy castings having boron added to improve transverse stress rupture strength and ductility of DS castings. The castings are heat treated at 1250° C. for 4 h so that a fully solution of the secondary phase (γ′-phase) is performed. Due to the occurrence of grain boundary cracks after the fully solution heat treatment the producibility is so deficient as to render DS columnar grain castings produced from the aforementioned nickel base superalloy unsuitable for use as turbine blades of gas turbine engines.

An object of the present invention is to provide a heat treatment of alloys, especially of as-cast alloys, e.g. DS columnar grain castings based on the aforementioned single crystal nickel base superalloy, having substantially improved transverse stress rupture strength and ductility as well as producibility to an extent that the DS castings are acceptable for use as high temperature applications such as turbine blades of a gas turbine engine.

SUMMARY OF THE INVENTION

The present invention involves a heat treatment of cast alloys, such as superalloys, having at least one addition, which improves grain boundary strength such as boron in the nickel base superalloy described here above in a manner discovered to significantly improve transverse stress rupture strength and ductility of directionally solidified (DS) columnar grain castings produced with a heat treatment, which solves a secondary phase only partly, e.g. no fully solution heat treatment is performed.

Boron is often added to superalloy compositions in an effective amount to substantially improve transverse stress rupture strength and ductility of directionally solidified columnar grain castings produced from the boron-modified superalloy. The boron concentration preferably is controlled in the range of about 0.003% to about 0.0175% by weight of the superalloy composition to this end.

In conjunction with addition of boron to the superalloy composition, the carbon concentration preferably is controlled in the range of about 0.05% to about 0.11% by weight of the superalloy composition.

A preferred nickel base superalloy in accordance with an embodiment of the present invention consists essentially of, in weight %, of about 11.6% to 12.70% Cr, about 8.50% to 9.5% Co, about 1.65% to 2.15% Mo, about 3.5% to 4.10% W, about 4.80% to 5.20% Ta, about 3.40% to 3.80% Al, about 3.9% to 4.25% Ti, about 0.05% to 0.11% C, about 0.003% to 0.0175% B, and balance essentially Ni. The boron modified nickel base superalloy can be cast as DS columnar grain castings pursuant to conventional DS casting techniques such as the well known Bridgman mould withdrawal technique.

DS castings produced in this manner typically have a plurality of columnar grains extending in the direction of the principal stress axis of the casting with the <001> crystal axis generally parallel to the principal stress axis. DS columnar grain castings pursuant to the present invention preferably exhibit a stress rupture life of at least about 100 hours and elongation of at least about 2.5% when tested at a temperature of 750 degrees C. (1382 degrees F.) and stress of 660 MPa (95.7 Ksi) and will find use as turbine blades, vanes, outer air seals and other components of a industrial and aero gas turbine engines. In many cases it is sufficient that the alloy after the heat treatment has mechanical properties which are less, particularly less than 10 or 20% compared to an alloy wherein the secondary phase is completely solved.

The above objects and advantages of the present invention will become more readily apparent form the following detailed description taken with the following drawings.

DETAILED DESCRIPTION OF THE INVENTION

Exemplarily as alloy a nickel base superalloy is chosen which consists essentially of, in weight %, of about 9.5% to 14% Cr, about 7% to 11% Co, about 1% to 2.5% Mo, about 3% to 6% W, about 1% to 6% Ta, about 3$ to 4% Al, about 3% to 5% Ti, about 0% to 1% Nb, and balance essentially Ni and B present in an amount effective to substantially improve transverse stress rupture strength of a DS casting as compared to a similar casting without boron present.

The inclusion of boron, as an addition, which improves the grain boundary strength in the alloy, is chosen in an amount discovered effective to provide substantial transverse stress rupture strength and ductility of a DS columnar grain casting produced from the alloy as compared to a similar casting without boron present.

Preferably, the nickel base superalloy is modified by the inclusion of boron B in the range of about 0.003% to about 0.0175%, preferably 0.010% to 0.015%, by weight of the superalloy composition to this end.

In conjunction with addition of boron to the superalloy composition, the carbon C concentration is controlled in a preferred range of about 0.05% to about 0.11% by weight of the superalloy composition. Also Silicon Si, Zirconium Zr and Hafnium Hf can be used as addition.

Furthermore all combinations of B, C, Si, Zr, Hf are possible. The addition of Zr or of Hf should not exceed 0.0075 wt % of the alloy.

The transverse stress rupture strength and ductility as well as the producibility of DS castings produced from nickel base superalloy with the modified heat treatment are provided to an extent that the castings are rendered acceptable for use as turbine blades and other components of gas turbine engines.

A particularly preferred boron-modified nickel base superalloy casting composition consists essentially of, in weight %, of about 11.6% to 12.70% Cr, about 8.5% to 9.5% Co, about 1.65% to 2.15% Mo, about 3.5% to 4.10% W, about 4.80% to 5.20% Ta, about 3.40 to 3.80% Al, about 3.9% to 4.25% Ti, about 0.05% to 0.11% C, about 0.003% to 0.0175% B, and balance essentially Ni and castable to provide a DS columnar grain microstructure.

The DS microstructure of the columnar grain casting typically includes about 0.4 to about 1.5 volume % of a phase based an tantalum carbide.

Although not wishing to be bound by any theory, it is thought that boron and carbon tend to migrate to the grain boundaries in the DS microstructure to add strength and ductility to the grain boundaries at high service temperatures, for example 816 degrees C. (1500 degrees F.) typical of gas turbine engine blades. DS columnar grain castings produced from the above boron modified nickel base superalloy typically have the <001> crystal axis parallel to the principal stress axis of the casting and exhibit a stress rupture life of at least about 100 hours and elongation of at least about 2.5% when tested at a temperature of 750 degrees C. (1382 degrees F.) and stress of 660 MPa (95.7 Ksi) applied perpendicular to the <001> crystal axis of the casting.

For example, the following DS casting tests were conducted and are offered to further illustrate, but not limit, the present invention.

A heat #1 having a nickel base superalloy composition in accordance with the aforementioned U.S. Pat. No. 4,597,809 and heats #1A and #2 and #3 of boron modified nickel base superalloy were prepared with the following compositions, in weight percentages, set forth in Table I: TABLE I Heat Cr Co Mo W Ta Al Ti C B Ni #1 12.1 9.0 1.8 3.7 5.2 3.6 4.0 0.07 0.001 balance #1A 12.1 9.0 1.8 3.7 5.2 3.6 4.0 0.08 0.010 balance #2 12.1 9.0 1.8 3.7 5.2 3.6 4.0 0.09 0.011 balance #3 12.1 9.0 1.8 3.7 5.2 3.6 4.0 0.08 0.014 balance

Each heat was cast to form DS columnar grain non-cored castings having a rectangular shape for transverse stress rupture testing pursuant to ASTM E-139 testing procedure. The DS castings were produced e.g. using the conventional Bridgman mould withdrawal directional solidification technique.

For example, each heat was melted in a crucible of a conventional casting furnace under a vacuum of 1 micron and superheated to 1427 degrees C. (2600 degrees F.). The superheated melt was poured into an investment casting mould having a face coat comprising zircon backed by additional slurry/stucco layers comprising zircon/alumina. The mould was preheated to 1482 degrees C. (2700 degrees F.) and mounted an a chill plate to effect unidirectional heat removal from the molten alloy in the mould. The melt-filled mould an the chill plate was withdrawn from the furnace into a solidification chamber of the casting furnace at a vacuum of 1 micron at a withdrawal rate of 6-16 inches per hour.

The DS columnar grain castings were cooled to room temperature under vacuum in the chamber, removed from the mould in conventional manner using a mechanical knock-out procedure, heat treated at a temperature and for a duration in such way, that the solution of a secondary phase in the matrix is only partly performed.

The nickel based superalloy has as a secondary phase the γ′-phase.

For a specimen (e.g. nickel based superalloy) with the composition of claim 25, the inventive heat treatment is performed after casting at 1213° C. for at least 1 h, which is not the solution temperature of a secondary phase (e.g. y′ phase) for this alloy.

Also the temperature of 1250° C. (called fully solution temperature), which is normally used for a fully solution treatment, can be used but only as long as the secondary phase is not completely solved in the matrix.

The not solubilized amount of the secondary phase in the matrix is smaller than 90, 70, 50 or 30 vol % according to the geometry and producibility after the heat treatment, because grain boundary cracks are avoided, in order to increase the yield rate of specimens and desired mechanical properties of the specimen.

The alloy can have a single crystal structure or only having grains along one direction.

Optionally an ageing heat treatment can be performed for this composition at 1080° C. for at least 2 h after this solution heat treatment. Optionally followed by a second ageing heat treatment at 870° C. for at least 12 h.

Especially the inventive heat treatment is used for hollow specimen, especially blades, vanes, or liners because cracks do appear more often in walls, especially in thin walls, than in massive specimens after the normally used heat treatment after casting.

The inventive heat treatment leads to an increased grain boundary strength during this heat treatment, so that the yield rate (components without cracks) after the heat treatment is increased.

Also the transverse stress rupture of the component as final product is increased during use of the component at working conditions, because grain boundary strength is increased. The inventive method yields also good results for massive components, e.g. of a gas turbine.

The castings were also analysed for chemistry, and machined to specimen configuration.

Stress rupture testing was conducted in air at a temperature of 750 degrees C. (1382 degrees F.) and stress of 660 MPa (95.7 Ksi) applied perpendicular to the <001> crystal axis of the specimens.

The results of stress rupture testing are set forth in TABLE II below where LIFE in hours (HRS) indicates the time to fracture of the specimen, ELONGATION is the specimen elongation to fracture, and RED OF AREA is the reduction of area of the specimens to fracture. The BASELINE data corresponds to test data for Heat #1, and the #1A, #2 and #3 data corresponds to test data for Heat #1A, #2 and #3, respectively. The BASELINE data represent an average of two stress rupture test specimens, while the #1A, #2 and #3 data represent a single stress rupture test specimen. TABLE II Elonga- Red of #of Temperature Stress Life tion Area Alloy Tests ° C. (° F.) MPa (KSI) (h) (%) (%) Baseline 2 750 (1382) 660 (95.7) 0 0 0 #2 1 750 (1382) 660 (95.7) 182 2.6 6.3 #3 1 750 (1382) 660 (95.7) 173 3.7 10.7 #1A 1 750 (1382) 660 (95.7) 275 3.1 4.7

It is apparent from TABLE II that the DS columnar grain specimens produced from heat #1 exhibited in effect essentially no (e.g. zero hours stress rupture life) transverse grain boundary strength when tested at a temperature of 750 degrees C. (1382 degrees F.) and stress of 660 MPa (95.7 Ksi). That is, the specimens failed immediately to provide an essentially zero stress rupture life. Moreover, the elongation and reduction of area data were essentially zero. These stress rupture properties are so deficient as to render the DS columnar grain castings produced from heat #1 unacceptable for use as turbine blades of gas turbine engines.

In contrast, TABLE II reveals that DS columnar grain specimens produced from heat #1A exhibited a stress rupture life of 275 hours, an elongation of 3.1%, and a reduction of area of 4.7 and specimens from heat #2 exhibited a stress rupture life of 182 hours, an elongation of 2.6%, and a reduction of area of 6.3% when tested at a temperature of 750 degrees C. (1382 degrees F.) and stress of 660 MPa (95.7 Ksi). These stress rupture properties of the invention represent an unexpected and surprising improvement over those of specimens produced from heat #1 and render DS columnar grain castings produced from heats #1A, #2 and #3 more suitable for use as turbine blades and other components of gas turbine engines.

The present invention can be applied not only to the foregoing described superalloys, but also to alloys which comprises a matrix and from about 0.4 to 1.5 vol % of a phase based on tantalum carbide consisting substantially of, in weight percent, about 10-14.5% (preferably 10-13.5%) chromium; 8-10% cobalt; 1.25-2.5% molybdenum; 3.25-4.25% tungsten; 4.5-6% tantalum; 3.25-4.5% aluminum; 3-5% (preferably 3-4.75%) titanium; 0.0025-0.025% boron; up to about 0.02% zirconium; 0.05-0.15% carbon; and having no intentional addition of niobium; no intentional addition of hafnium; and balance essentially nickel; wherein aluminum+titanium is between about 6.5-8%. The alloy also includes roughly about 0.4 to 1.5 vol. % of a phase based on tantalum carbide. More preferably, the alloy comprises about 11-13% chromium; 8.25-9.75% cobalt; 1.5-2.25% molybdenum; 3.4-4.3 % tungsten; 4.7-5.5% tantalum; 3.3-4% aluminum; 3.75-4.3 titanium; 0.008-0.025% boron; up to about 0.02% zirconium; 0.08-0.13 carbon; wherein aluminum+titanium is between about 7-8%. Most preferably, the alloy comprises about 12% chromium; 9% cobalt; 1.9% molybdenum; 3.8% tungsten; 5% tantalum; 3.6% aluminum; 4.1% titanium; 0.015% boron; less than 0.02% zirconium; 0.10% carbon; and having no intentional addition of zirconium (and in any event less than about 0.02 Zr) and no intentional addition of niobium; no intentional addition of hafnium; balance essentially nickel.

The following tables III to VI shows some embodiments of compositions to be used. TABLE III Cr Ti Mo W Ta Al Co B Zr C Hf #4 11.6 4.0 1.8 3.8 5.1 3.6 8.9 0.005 0.014 0.07 0.48 #5 11.7 4.0 1.8 3.7 5.0 3.6 8.9 0.005 0.016 0.06 0.86 #6 12.3 4.0 1.8 3.7 5.0 3.5 8.8 0.018 0.092 0.11 0.47 #7 11.9 4.0 1.8 3.8 5.2 3.5 8.9 0.008 0.021 0.06 0.01 #8 11.6 4.0 1.8 3.7 5.3 3.6 8.9 0.008 0.031 0.07 0.50 #9 11.9 4.0 1.8 3.7 4.9 3.5 8.8 0.019 0.101 0.12 0.92

TABLE IV Cr Ti Mo W Ta Al Co B Zr C Hf #10 11.6 4.0 1.8 3.8 5.1 3.6 8.9 0.005 0.014 0.07 0.0070 #11 11.7 4.0 1.8 3.7 5.0 3.6 8.9 0.005 0.015 0.06 0.0060 #12 12.3 4.0 1.8 3.7 5.0 3.5 8.8 0.018 0.091 0.11 0.0075 #13 11.9 4.0 1.8 3.8 5.2 3.5 8.9 0.008 0.020 0.06 0.0072 #14 11.6 4.0 1.8 3.7 5.3 3.6 8.9 0.008 0.032 0.07 0.0066 #15 11.9 4.0 1.8 3.7 4.9 3.5 8.8 0.019 0.103 0.12 0.0054

TABLE V Cr Ti Mo W Ta Al Co B Zr C Hf #16 11.6 4.0 1.8 3.8 5.1 3.6 8.9 0.005 0.0065 0.07 0.0072 #17 11.7 4.0 1.8 3.7 5.0 3.6 8.9 0.005 0.0055 0.06 0.0062 #18 12.3 4.0 1.8 3.7 5.0 3.5 8.8 0.018 0.0072 0.11 0.0073 #19 11.9 4.0 1.8 3.8 5.2 3.5 8.9 0.008 0.0075 0.06 0.0071 #20 11.6 4.0 1.8 3.7 5.3 3.6 8.9 0.008 0.0070 0.07 0.0064 #21 11.9 4.0 1.8 3.7 4.9 3.5 8.8 0.019 0.0048 0.12 0.0058

TABLE VI Cr Ti Mo W Ta Al Co B Zr C Hf #22 11.6 4.0 1.8 3.8 5.1 3.6 8.9 0.005 0.0042 0.07 0.48 #23 11.7 4.0 1.8 3.7 5.0 3.6 8.9 0.005 0.0057 0.06 0.86 #24 12.3 4.0 1.8 3.7 5.0 3.5 8.8 0.018 0.0056 0.11 0.47 #25 11.9 4.0 1.8 3.8 5.2 3.5 8.9 0.008 0.0063 0.06 0.01 #26 11.6 4.0 1.8 3.7 5.3 3.6 8.9 0.008 0.0038 0.07 0.50 #27 11.9 4.0 1.8 3.7 4.9 3.5 8.8 0.019 0.0075 0.12 0.92

The present invention can also be applied to a nickel based superalloys consisting essentially of, in weight %, 12.5 to about 15% Cr, greater than 5% to less than 15% Co, 2.5% to 5% Mo, 3% to 6% W, 2% to 4% Al, 4% to 6% Ti, 0.005% to 0.02% B, up to 0.1% Zr, and balance essentially nickel and carbon with the ratio of Ti to Al being greater than 1 but less than 3, the sum of Ti and Al being 7.5-9 weight %; the sum of Mo and half of the W being 5-7 weight %; and with carbon content maintained below 0.08% to unexpectedly improve machinability after appropriate heat treatment such as solution heat treatment and precipitation hardening heat treatment steps by virtue of beneficially affecting primary carbides in the alloy microstructure, while providing acceptable mechanical properties. The Cr concentration preferably is reduced in the range of about 13 to about 14 weight %, preferably nominally 13.5 weight % Cr, to compensate for the lower carbon content of the alloy of the invention.

A more narrow composed nickel base superalloy consists essentially of, in weight %, of about 12.5% to 15% Cr, about 9.00% to 10.00% Co, about 3.70% to 4.30% Mo, about 3.70% to 4.30% W, about 2.80% to 3.20% Al, about 4.80% to 5.20% Ti, about 0.005% to 0.02% B, up to about 0.10% Zr, less than about 0.08% C, and balance essentially Ni. Preferably, the nickel base superalloy is modified by reducing carbon in the range of abut 0.055% to about 0.075% by weight, preferably about 0.07% by weight, of the superalloy composition to this end.

A particularly preferred carbon modified nickel base superalloy casting composition consists essentially of, in weight %, nominally about 13.50% Cr, about 9.40% Co, about 4.0% Mo, about 4.00% W, about 3.00% Al, about 5.00% Ti, about 0.015% B, about 0.07% C, and balance essentially Ni.

One preferable nickel based superalloy is modified by reducing carbon in the range of below about 0.08 weight %, wherein, in weight %, the nickel base alloy consists essentially of about 12.5 to 15 % Cr, about 9.00% to 10% Co, about 3.70% to 4.30% Mo, about 3.70% to 4.3% W, about 2.80% to 3.20% Al, about 4.80% to 5.20% Ti, about 0.005% to 0.02% B, up to about 0.10% Zr, and balance essentially Ni.

The present invention is effective to provide DS columnar grain castings with substantial transverse stress rupture strength and ductility. These properties are achieved without adversely affecting other mechanical properties, substantially such as tensile strength, creep strength, fatigue strength, and corrosion resistance of the DS castings. The present invention is especially useful to provide large DS columnar grain industrial gas turbine (IGT) blade castings which have the alloy composition described above to impart substantial transverse stress rupture strength and ductility to the castings and which have a length of about 20 centimeters to about 60 centimeters and above, such as about 90 centimeters length, used throughout the stages of the turbine of stationary industrial gas turbine engines. The above described boron-modified nickel base superalloy casting composition can be cast as DS columnar grain or single crystal components.

While the invention has been described in terms of specific embodiments thereof, it is not intended to be limited thereto but rather only to the extent set forth in the following claims. 

1. A method of heat treating a directionally solidified columnar grain nickel base alloy casting to improve grain boundary strength, comprising casting the alloy such that it has a secondary phase after casting which is only partially solved in the matrix of the alloy at a solution temperature, the alloy comprising Boron present in an amount effective to substantially improve transverse stress rupture strength of the casting.
 2. A method of claim 1, wherein B is present in the range of about 0.003% to about 0.018% by weight.
 3. A method of claim 1, wherein after the heat treatment the alloy has a stress rupture life of at least about 100 hours and elongation to fracture of at least about 2.5% when tested at a temperature of 750 degrees C. (1382 degrees F.) and stress of 660 MPa (95.7 Ksi) applied in a direction perpendicular to a <001> crystal axis of the casting.
 4. A method of claim 1, wherein the alloy consists essentially of, in weight %: about 11.6% to 12.70% Cr, about 8.5% to 9.5% Co, about 1.65% to 2.15% Mo, about 3.5% to 4.10% W, about 4.8% to 5.20% Ta, about 3.4% to 3.80% Al, about 3.9% to 4.25% Ti, about 0.05% to 0.11% C, about 0.003% to 0.0175% B, balance essentially Ni and having substantially improved transverse stress rupture strength as compared to a similar casting without boron present.
 5. A method of claim 1, wherein the alloy consists essentially of, in weight %: about 12.00% Cr, about 9.00% Co, about 1.85% Mo, about 3.70% W, about 5.10% Ta, about 3.60% Al, about 4.00% Ti, about 0.0125% B, about 0.09% C, balance essentially Ni and having a stress rupture life of at least about 100 hours and elongation to fracture of at least about 2.5% when tested at a temperature of 750 degrees C. (1382 degrees F.) and stress of 660 MPa (95.7 Ksi) applied perpendicular to a <001> crystal axis of said casting.
 6. The method of claim 1, wherein the heat treatment is performed after casting.
 7. The method of claim 6, wherein a fully solution temperature is used.
 8. The method of claim 1, wherein the alloy is formed into a hollow component selected from the group consisting of vanes, blades and liners.
 9. The method of claim 1, wherein the alloy further comprises an addition selected from the group consisting of Zirconium, Silicon, Hafnium.
 10. The method of claim 9, wherein the Zirconium does not exceed 0.0075wt % of the alloy.
 11. The method of claim 9, wherein the Hafnium does not exceed 0.0075wt % of the alloy. 